KR20230131181A - New powder, method for additive manufacturing of parts made from the new powder, and articles made therefrom - Google Patents

Abstract

The present invention provides Al-based powders suitable for additive manufacturing. In one embodiment, the powder contains 3% to 5.5% Mn, 0.2% to 2% Zr, 0.1% to 1.4% Cr, 0% to 2% Mg, totaling up to 0.7%. It consists of Fe and Si in weight percent, O as an inevitable impurity up to 0.7 weight percent, and other inevitable impurities up to 0.5 weight percent, and the balance Al. The invention also relates to additive manufacturing methods as well as articles manufactured by additive manufacturing.

Description

New powder, method for additive manufacturing of parts made from the new powder, and articles made therefrom

The present invention relates to a novel aluminum alloy suitable for use in additive manufacturing (AM) of components made from the alloy.

Metal-based additive manufacturing has applications in many industries, including the aerospace and automotive industries.

Aluminum alloys are useful in a variety of applications. Aluminum alloy products are generally manufactured through shape casting or forging processes. Shape casting generally involves casting a molten aluminum alloy into its final form, such as through high pressure dies, permanent molds, green and dry sand, investment, or plaster casting. Forged products are generally manufactured by casting molten aluminum alloy into ingots or billets. The ingot or billet is generally further hot worked, and sometimes cold worked, to produce its final form.

Aluminum alloys can be divided into several systems depending on alloy content and intended processing. Most Al-based casting alloys contain 9 to 12% Si, the element forming a eutectic at about 12% Si, which has excellent casting properties. Other elements that can be used in casting alloys include Mg and Cu. The addition of these elements maintains excellent casting properties combined with strengthening by solid solution and precipitation of intermetallic compounds. Typical examples include Al-7Si-0.7Mg and A19Si-3Cu.

To reach higher strengths, it is necessary to use other alloying elements and reduce the content of Si. Forged aluminum alloys with higher strengths are named 2000-, 6000-, and 7000-series alloys. Tables 1 and 2 show the chemical composition and mechanical properties of common aluminum alloys.

Table 1.

Al-Cu alloy (2000-series)

In this group of alloys Mg is a secondary addition. Other elements include Mn, Cr, and Zr. The main strength contribution is due to solution hardening and precipitation hardening by intermetallic phases, for example Al ₂₀ Cu ₂ Mn ₃ , Al ₁₈ Mg ₃ Cr ₂ or AlZr ₃ . To achieve this hardening, a solid solution heat treatment is performed, followed by an aging treatment. Iron and silicon are considered impurities because they form harmful intermetallic compounds. Titanium may be added to control grain structure during ingot casting. The corrosion properties of this group of alloys are less favorable compared to other types of Al-alloys. Parts made from these alloys have a high strength-to-weight ratio and can be used up to 150°C. The alloy generally cannot be welded but has excellent machinability.

Al-Mg-Si alloy (6000-series)

This group of alloys is the work horse of high-strength Al-alloys. Mg and Si additions are generally carried out in the ratio of Mg ₂ Si as the main precipitate. The alloy will require solution heat treatment and aging but will not reach the strength levels of the 2000- and 7000-series alloys. However, they are weldable and exhibit excellent formability, extrudability, machinability and corrosion resistance.

Al-Zn alloy (7000-series)

This group of alloys has Zn as the main alloying element with additions of Mg, Cu and Cr. These elements form intermetallic phases used in solution heat and aging treatments at very high strengths. Resistance to stress corrosion cracking is reduced, and these materials are often used under overaging conditions to reach the best combination of strength, fracture toughness and corrosion resistance.

The mechanical properties of different groups of alloys are summarized in Table 2.

Table 2.

Corrosion is a system property in that it depends on the material as well as the environment in which it operates. Nevertheless, it is possible to make some general predictions about the behavior of alloys. Aluminum typically forms a passive oxide layer that minimizes the general corrosion rate. If the layer is broken by local changes in composition, for example in the form of precipitates or interfaces between dissimilar metals, corrosion may occur in the form of crevice corrosion, pitting corrosion or galvanic corrosion. These are all processes that can lead to very rapid and severe destruction. Aluminum alloys in which the alloying elements are present in solution in the fcc-matrix generally show high resistance to all localized corrosion forms, whereas alloys with precipitates containing Cu or Zn perform poorly in this regard.

In a laser powder bed fusion (LPBF) process, a laser beam is scanned over a powder bed, melting the powder particles in a predetermined pattern, then a new powder layer is applied, and the process is repeated. To manufacture a fully dense part, it is necessary to melt not only the powder particles of the new layer, but also some of the previously formed layers. By doing so, complete fusion between the layers can be ensured and any defects such as porosity can be eliminated. The laser scans very quickly (200 to 7000 mm/s) and the laser spot is small (40 to 100 μm); This means that the melt pool formed is very small and the time the material is in the liquid state is very short (about 0.2 msec). These manufacturing limits aim to ensure that any separation of alloying elements during solidification is minimized due to the short interaction time available.

However, the directional nature of the process causes both grain and subgrain structures to be oriented towards the heat source. The subcrystal grain shape is dendritic or cellular. The structure is typically large columnar grains with very fine cell structures in the building direction. Most of the heat extracted from the manufactured part escapes through the part to the build plate, which acts as a heat sink, enhancing the directional nature of the solidification process.

The traditional alloys described above are not always suitable for the LPBF process. In general, it is only cast alloys that can be adopted for LPBF without problems. A360.0 is a reasonable process window suitable for alloys and industrial processes that print very well with few defects. However, this material provides only limited strength.

Alloys that have reached sufficient strength using traditional manufacturing processes are much more difficult to machine without significant defects using the LPBF process. In particular, the solidification behavior is unfavorable, leading to cracking during solidification.

Several inventors have partially addressed these problems:

EP0105595A2 (Alcan) discloses an aluminum alloy consisting of 1.5 to 7.0 wt.% of Cr, 0.5 to 2.5 wt.% of Zr, 0.25 to 4.0 wt.% of Mn, and the remainder of aluminum and general impurities.

In another embodiment, the composition ranges from 3.0 to 5.5 wt.% Cr, 1.0 to 2.0 wt.% Zr, 0.8 to 2.0 wt.% Mn, the balance Al and general impurities.

The inventors also disclose a method of making powders by rapidly solidifying the molten alloy at a cooling rate of at least 1000° C./sec and fast enough to produce soft particulates in which the bulk of the alloy additives remains in solid solution.

The particles thus formed are solidified at a temperature of 300 to 500° C. to achieve age hardening. In one example, the powder of Alloy A (Cr 5.25, Zr 1.75, Mn 1.75, remainder Al) is produced by gas atomization and then by sheet rolling at 350°C, which leaves areas without visible deposits and areas with deposits. Creates a mixed microstructure with zones. The tensile strength of the material is 588 MPa, the yield strength is 530 MPa, and the elongation at break is 6%.

CN 11659889 A has Mn: 1.00 wt.% to 10.00 wt.%, Mg: 0.01 wt.% to 3.00 wt.%, Si: 0.01 wt.% to 2.00 wt.%, Zr: 0.01 wt.% to 3.50 wt. %, Fe: 0.01 wt.% to 1.50 wt.%, the remainder being Al. A high-strength aluminum-manganese alloy is disclosed.

WO2020/139427 (HRL Laboratories) discloses additively manufactured high temperature aluminum alloys and feedstocks for producing them. The disclosed feedstock consists of a powder mixture comprising 80 to 99 weight percent of an aluminum-containing base powder mixed with 1 to 20 weight percent of an alloyed powder having a particle size smaller than the aluminum-containing base powder. In certain embodiments, the feedstock contains additional alloying elements.

In one example, a gas atomized base powder is disclosed containing the following elements in weight percent: Al 92, 6, Cu 6.7, Mn 0.35, Ti 0.24. A base powder having a particle size distribution suitable for selective laser melting; D10=15 microns, D50=27 microns and D90=44 microns. The base powder is mixed with zirconium powder having an average particle size of 0.5 to 1.5 microns at a content of 2% by weight.

WO2018/009359A1 (NanoAl) discloses an aluminum alloy family with high strength and ductility, excellent corrosion resistance and weldability.

Disclosed herein are aluminum alloys comprising about 1 to 10 weight percent Mg, 0.45 to 3 weight percent Zr, and the balance Al. The alloy is completely free of intentionally added scandium, erbium, thulium, ytterbium or lutetium.

In another embodiment, an aluminum alloy comprising Mg and at least one element from group 4 elements Ti, Zr and Hf, group 5B elements V, Nb, Ta, group 6B elements Cr, Mo and W, the remainder Al is disclosed. It is done. An alloy completely free of intentionally added scandium, erbium, thulium, ytterbium, or lutetium, wherein the alloy contains nano-scale aluminum-transition ion metal precipitates in an aluminum matrix having an average diameter ranging from about 3 nm to about 50 nm. and the transition metal is selected from the group of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo and W.

Scalmalloy/Al-Mg-Sc alloy was recently developed for laser powder bed fusion process.

EP3165620A1 and EP3181711A1 (Airbus Defense and Space GmbH) teach that high strength, elongation and hardness can be achieved by laser powder bed fusion processing. The manufactured parts are defect-free and exhibit high corrosion resistance. The invention is based on the use of scandium, an expensive and rare rare earth element.

Additive manufacturing (AM) of objects formed from powdered aluminum-alloys is limited today by the lack of alloy powders suitable for manufacturing high-strength objects. Most available high-strength Al-alloy compositions have been developed for casting and used in the form of powders for AM, for example, by laser powder bed fusion, and the formed objects are prone to cracking. Some newly developed aluminum alloys used in laser powder bed fusion are believed to enable the production of crack-free objects.

There is therefore a need for an aluminum powder alloy that can be produced from a homogeneous melt, for example by spraying, without forming excessive precipitates that interfere with the manufacturing process. Additionally, these powders would be suitable for forming crack-free objects with high strength through AM processes such as LPBF, both as printed and after heat treatment.

The object of the present invention is to provide an aluminum powder suitable for the additive manufacturing of crack-free articles, in particular for use in laser powder bed fusion, LPBF. The alloy composition can be prepared entirely as a pre-alloy powder, avoiding separation of particles and non-uniform powder composition, and can be free of expensive alloying rare earth elements and copper. In particular, the alloy of the present invention will be completely free of intentionally added rare earth metals such as scandium, erbium, thulium, ytterbium or lutetium and will not contain intentionally added Cu. This object is met in a first aspect of the invention.

Another object of the present invention is to provide articles made from powder, which are free of cracks and have high strength both as printed and after heat treatment. This object is met in the second aspect of the invention.

A further object of the present invention is to provide a method for additive manufacturing of articles made from powder. The provided method is robust and has a so-called wide process window, i.e. variations in process parameters that do not affect the properties of the printed object. This object is met in the third aspect of the invention.

Additionally, because chromium is an environmental contaminant, the use of high levels of chromium in consumer products in general and in particular is undesirable. Accordingly, an aspect of the present invention is to provide a low chromium Al-based alloy that has the above desired properties or provides the above desired properties when used.

The above-mentioned object is achieved by the aspects and embodiments thereof as set forth in the appended claims. Aspects of the invention are summarized below under the heading Summary of the invention.

Summary of the invention

In a first aspect and embodiment, described in detail herein are pre-alloyed Al-based powders suitable for additive manufacturing consisting of the following total pre-alloyed Al-based powders by weight:

3 to 5.5% by weight Mn,

0.2 to 2% by weight Zr,

0.2 to 1.4% by weight Or,

0 to 2% by weight Mg,

Fe and Si totaling up to 0.7% by weight,

0 as unavoidable impurities up to 0.7% by weight, and additional unavoidable impurities up to 0.5% by weight,

The rest of Al.

In an embodiment thereof, a pre-alloyed Al-based powder according to claim 1, wherein the pre-alloyed Al-based powder has a Zr content of 0.3 to 1.8% by weight, is described.

In an embodiment thereof, a pre-alloyed Al-based powder is described, with a Cr content of 0.3 to 1.4% by weight.

In an embodiment thereof, a pre-alloyed Al-based powder is described, with a total content of Fe and Si of at most 0.2% by weight.

In an embodiment thereof, a pre-alloyed Al-based powder is described, wherein the pre-alloyed Al-based powder does not contain intentionally added Mg.

In an embodiment thereof, a pre-alloyed Al-based powder is described having a solidification cracking susceptibility index S according to equation (1) below 20°C.

In an embodiment thereof, a pre-alloyed Al-based powder is described, having a particle size range of 10 to 53 μm as measured by laser diffraction according to ISO 13320-1 1999, where less than 10% of the particles are outside the specified particle size range. .

In a second aspect, a process for manufacturing an article by additive manufacturing comprising the following steps is described in detail herein:

- providing a pre-alloyed Al-based powder according to any one of the embodiments of the first aspect;

- depositing a layer of pre-alloyed Al-based powder to form a powder layer;

- heating the pre-alloyed Al-based powder according to a predetermined pattern by laser or electron-beam to fuse or melt the particles to be bonded;

- cooling the fused or molten particles;

- repeating the deposition, heating and cooling steps until an additively manufactured article is formed;

- Recovering manufactured goods.

In this embodiment, the manufactured article subsequently undergoes heat treatment.

In an embodiment thereof, the heat treatment comprises heating the manufactured article to a temperature of 150 to 450° C. for a period of at least 0.5 hours in an air atmosphere.

In a third aspect, an article formed by a melt fusion process of pre-alloyed Al-powder according to any embodiment of the first aspect is described in detail.

In an embodiment thereof, an article is described, which is formed by a thermal fusion process of pre-alloyed Al-powder, wherein the melt fusion process is according to any of the embodiments of the second aspect.

Figure 1.
Simulation of coagulation according to the Scheil-Gulliver law. The line represents the temperature at which the remaining liquid (solidus) is in equilibrium with the solid as a function of the mass fraction solid formed.
Figure 2.
Process window diagram for alloy D. The relative density is shown as a contour diagram with laser power (W) and laser speed (mm/s) on the ordinate and abscissa axes, respectively. The hatch distance is 100 μm. The process window is the area where the relative density is greater than 98.5%.
Figure 3.
Hardness as a function of heat treatment. The aging treatment was carried out at 410°C in air. Time at temperature is shown on the abscissa.
Figure 4:
Investigation of cracking susceptibility to combined Fe and Si for alloys of the present invention.
Figure 5:
(a) Particle size (μm) versus volume %, and
(b) Particle size (μm) versus cumulative volume %.
Figure 6:
DOE cube design showing notches along all three sides and front side to indicate gas flow direction (depending on build direction and gas flow direction; front: XZ; top side: XY; right side: YZ)
Figure 7:
Design of experiments (DOE1) representing different processing parameters and number of samples per parameter. Layer thickness and laser power were kept constant.
Figure 8:
Design of Experiment 6.1 Results showing surface plots of the relative densities of all four alloys, a) Alloy A b) Alloy B c) Alloy C d) Alloy D.
Figure 9:
Surface plots showing experimental designs (1, 2, 3) showing laser speed vs. hatch and laser speed vs. laser power for a) Experimental Design 6.2, b) Experimental Design 6.3, c)/d) Alloy D, respectively.
Figure 10:
a) X-ray plots for all four alloys in as-printed conditions. b) Plot for alloy D in as-printed conditions, with inset showing the region between 20 = 37°-47°, showing the minor phases formed in all four alloys. It is assumed to be an Al-Mn based precipitate. △ = pure-Al peak and ◆ = A1-Mn precipitate peak.
Figure 11:
a) Theoretical effects on the lattice parameters of Mn, Cr, and Zr. It is noted that the dashed line represents the intercept correction assuming pure Al = 4.0478 Å instead of Al = 4.016 Å.
b) Lattice parameter values based on X-ray diffraction patterns for all four alloys in as-printed conditions. Reasonable values within 0.01 Å were achieved.
Figure 12:
a) 20X optical microscope image of as-sprayed powder Alloy C, etched with Keller's reagent.
b) 50X optical microscopy image of as-sprayed powder Alloy C, etched with Keller's reagent.
Figure 13:
a) Top side (XY), b) Right side (YZ) and c) Front (XZ) views of a cube of alloy D as printed with -99.6% relative density.
Figure 14:
Two types of typical defects observed after printing of Al-alloys (exemplified for alloy D); a) shows a keyhole defect at low magnification (a) and magnification (b); Low magnification and magnification show lack of fusion defects (d).
Figure 15:
SEM image of as-sprayed powder of Alloy C showing dendritic structure (a); Enlarged image (b) with nanometer particles inside the dendritic region (indicated by yellow arrows). The particles indicated by the red arrows are silica left over from polishing with OP-U.
Figure 16:
a) SEM image of Alloy D (as printed) along the front side in bottom to top build direction and b) melt pool shown in red along the top side in inner build direction from the image.
Figure 17:
SEM image showing high resolution of melt pool boundary deposits. Two categories of deposits were identified: large (200 to 300 nm, blue arrows) and small (<100 nm, yellow arrows). Large deposits are shown at higher magnification in the inset in b).
Figure 18:
SEM image of Alloy D showing a) the melt pool boundary of the grain refinements in the melt pool boundary region and b) the enlarged region showing grain boundaries and subgrain boundary particles.
Figure 19:
Chemical composition of three categories of Mn-containing deposits shown by EDS analysis, including large and small deposits in subgrain boundary deposits as well as melt pool boundary deposits. Sub-grain boundary deposits are shown in Figure 18 and melt pool boundary deposits are visualized in Figure 17.
Figure 20:
SEM image of Alloy D showing two categories of melt pools: a) normal melt pools with long columnar grains epitaxial to the build direction, b) melt-pools with refined grains (grain refinement is shown in the inset image).
Figure 21:
a) SEM image in SE2 mode showing alloy D in the top face section at 20kX with Zr dendritic structure.
b) SEM image in BSE mode showing alloy C in the top face section at 2.5kX with Zr tetrahedral structure.

details

To optimize alloys for rapid solidification processes, Integrated Computational Materials Engineering (ICME) is powerful.

The effect of different alloying elements on the solidification behavior and precipitation process of the alloy can be calculated and good predictions of properties for new alloying systems can be obtained.

The method is based on an in-depth understanding of the relationships between phase relationships and properties for specifically targeted materials and processes. A combination of selected experimental studies and calculations is required.

The advantage of the ICME development process is that it allows investigation of previously unknown areas of the composition space for a class of materials. This method was used in the present invention to show why the alloy composition is fundamentally different from available alloys in terms of how alloying elements are used to achieve high strength components.

Alloying elements were selected based on their potential solid solution strengthening effects, as described through first principles calculations.

This was then combined with previous knowledge of solid solution expansion in Al-based alloys to estimate how much solute could be expected to dissolve and what the strength of the printed state of these alloys would be.

The amount of supersaturation in solid solution is supported by nonequilibrium calculations and evaluation of the resistance to nucleation/coarsening of the second phase particles based on diffusivity and rapid cooling during LPBF processing.

A tentative phase diagram (from binary to quaternary) is developed using ThermoCalc software (versions and databases), and possible subcooling through LPBF processing is used as a factor to determine how high the percentage of solute can be dissolved compared to equilibrium. It was used.

The second important criterion for alloy design was to select alloying elements based on Scheil solidification calculations and to select for negligible solidification cracking susceptibility.

This approach will create an easy path to avoid unnecessary problems during printing. Scheil curves were developed using ThermoCalc software using the TCAL7 and M0BAL5 databases.

Unexpectedly, it has been found that no low melting phase is formed near the end of the solidification stage for the alloys of the present invention, thus avoiding the risk of solidification cracking. This is indicated by the excellent crack solidification index.

Additionally, Cr-containing phases were previously expected to form during solidification in printing, reducing the solution hardening effect of Cr, and by careful selection of the element and its content, surprisingly, the formation of these precipitates could be prevented. Similarly, alloying elements Zr and Mn with contents according to the invention will not form harmful precipitates during printing and will enable age hardening.

The pre-alloyed Al-based powder was designed using Integrated Computational Materials Engineering (ICME) for additive manufacturing processes where solidification is fast, i.e., >1000° C./s, and the formation of precipitates in the liquid is suppressed.

Alloying elements play a role in strengthening the alloy and will only be effective in rapidly solidified microstructures. In all embodiments of the invention, the alloys have a stable solidification path up to 80% mole fraction of solids with a solidification temperature drop of less than 25°C, which is in contrast to existing prior art alloys (see Figure 1) except A360.0. This enables LPBF-printing (see Figure 4) with essentially no crack formation (see Figure 1). Additionally, the most preferred alloys of the present invention have a stable solidification path, i.e., upon solidification, a single structure is formed and can be printed, for example, by LPBF without cracking (see Figure 4).

The article in the printed state is characterized by a structure characterized by a cellular/dendritic microstructure with a single-phase phase composition of the microstructure, containing less than 5% by volume of Al ₆ Mn, Al ₃ Zr and other precipitates.

The pre-alloyed Al-alloyed powder according to the invention can be used in the additive manufacturing of Al-alloy parts. Parts can be printed easily with a wide processing window. The parts exhibit high strength and can be heat treated after the additive manufacturing step to further improve properties.

Accordingly, to obtain improved strength, the article in the printed state can be heat treated, i.e. age-hardened at 150 to 450° C. for a period of at least 0.5 hours in an air atmosphere. Thereby, articles can be produced having a hardness of at least 125 HVO.3, measured according to ISO 6507-1 - 2018.

The microstructure of the article is characterized by a cellular/dendritic microstructure, with the phase composition of the microstructure being a single phase containing 2 to 15% by volume of Al ₆ Mn and Al ₃ Zr, and up to 10% by volume of other precipitates.

The alloying system is based on Al-Mn with additions of Cr and Zr. The results from experiments and calculations are particularly presented in Table 3, where each component of the aluminum alloy is denoted with an " Contributes to the effect.

Compositional areas of particular interest are highlighted in Table 3, where the elements of this alloy exhibit optimal synergistic effects for this alloy. However, although it is less desirable to work in the concentration region surrounding this optimal region of interest, alloys of these sub-optimal elemental inclusions nevertheless exhibit improved properties compared to other alloys of the prior art and are therefore excluded from the present invention. It doesn't work.

According to the invention, the following is described in detail herein:

Pre-alloyed Al-based powders suitable for additive manufacturing consisting of the following total pre-alloyed Al-based powders by weight:

3 to 5.5% by weight Mn,

0.2 to 2% by weight Zr,

0.2 to 1.4% Cr by weight,

0 to 2% by weight Mg,

Fe and Si totaling up to 0.7% by weight,

0 as unavoidable impurities up to 0.7% by weight, and additional unavoidable impurities up to 0.5% by weight,

The rest of Al.

Oxygen is generally present ubiquitously in aluminum alloys as an unavoidable impurity. The inventive process for making the inventive pre-alloy powder provides 0.35 to 0.55% by weight of oxygen as an inevitable impurity (see Table 9), but depending on the powder size, this value will vary. No specific or deleterious effects of oxygen in the specified ranges were observed in the experiments.

In general, it is desirable for the level of additional unavoidable impurities to be as low as technically possible. Preferably, the concentration of additional unavoidable impurities is less than 0.4 wt.%, less than 0.3 wt.% or even more preferably less than 0.2 wt.% or less than 0.1 wt.% of the pre-alloyed Al-based powder.

According to the invention, the alloying elements Cr and Mn are provided to bring about a solution hardening effect, while the alloying elements Zr are provided to contribute to precipitation hardening during subsequent heat treatment after the printing process. Surprisingly, as shown in the experiments and highlighted in Table 3, a synergistic effect of the added components was observed in the claimed concentration range.

The content of Zr is 0.1 to 2.0% by weight of the pre-alloyed Al-based powder. Experiments (see Table 3) have shown that Zr-contents above 2% will interfere with the spraying process and contents below 0.1% by weight will not contribute to the desired precipitation hardening effect after heat treatment of the article.

Preferably, the zirconium content is at least 0.2 wt.% of the pre-alloyed Al-based powder, at least 0.3 wt.% of the pre-alloyed Al-based powder and in articles made from the pre-alloy Al-based powder, at least 0.4 wt. .% or more, more preferably 0.5 wt.% or more, or even more preferably 0.6 wt.% or more.

Preferably, the zirconium content is less than 2% by weight of the pre-alloyed Al-based powder, less than 1.9 wt.% of the pre-alloyed Al-based powder and articles made from the pre-alloyed Al-based powder, less than 1.8 wt.%. %, less than 1.7 wt.%, more preferably less than 1.6 wt.%, less than 1.5 wt.%, less than 1.4 wt.%, less than 1.3 wt.%, or less than 1.25 wt.%.

In a preferred embodiment, the zirconium content in the pre-alloyed Al-based powder and in the articles made from the pre-alloyed Al-based powder is 0.2 to 1.9 wt.%, preferably 0.3 to 1.8 wt.%, 0.4 to 1.7 wt. .%, more preferably 0.5 to 1.6 wt.%, 0.5 to 1.5 wt.%, most preferably 0.6 to 1.4 wt.%, 0.7 to 1.3 wt.%, or 0.8 to 1.1 wt.%.

According to the present invention, it was considered that the addition of Mn and Cr would result in increased strength due to substitution of Al atoms for Mn and Cr atoms. If the matrix is supersaturated by Mn and Cr, precipitates will form in the fcc-Al matrix, which will lead to a decrease in the substitution strengthening effect.

Surprisingly (see Table 3), for a given concentration of Cr and Mn, the presence of both Cr and Mn in the alloy of the present invention was found to be significantly lower than expected crack formation in articles formed by LPBF-printing. Based on general principles, it was considered that this effect would be stronger with higher additions of Mn and Cr, but surprisingly, the effect peaks at about 0.7 wt.% Cr, with the chromium content increasing to 1.4 wt.% in the alloy of the present invention. It has been found to disappear when replaced.

In subsequent experiments, it was found that addition of Cr content exceeding 1.4% to the compositions of the present invention would result in the formation of undesirable intermetallic Al _x Cr _y precipitates in the melt at temperatures above 800°C.

These precipitates that form in the melt are so large that they do not contribute to strengthening the alloy and will remove Cr from the solution reducing substitutional strengthening. Additionally, the formation of intermetallic phases at high temperatures causes problems in spraying along with clogging of the nozzle.

The content of Cr is 0.1 to 1.4% by weight of the pre-alloyed Al-based powder. Preferably, the chromium content is at least 0.2 wt.% of the pre-alloyed Al-based powder, at least 0.3 wt.% of the pre-alloyed Al-based powder and in articles made from the pre-alloy Al-based powder, at least 0.4 wt. .% or more, more preferably 0.5 wt.% or more, or even more preferably 0.6 wt.% or more.

Preferably, the chromium content is less than 1.4 wt.% of the pre-alloyed Al-based powder, less than 1.35 wt.% of the pre-alloyed Al-based powder and articles made from the pre-alloy Al-based powder, less than 1.30 wt. %, less than 1.25 wt.%, more preferably less than 1.20 wt.%, less than 1.15 wt.%, less than 1.10 wt.%, less than 1.05 wt.%, or less than 1.0 wt.%.

In a preferred embodiment, the chromium content in the pre-alloyed Al-based powder and in articles made from the pre-alloyed Al-based powder is 0.2 to 1.4 wt.%, preferably 0.3 to 1.35 wt.%, 0.4 to 1.30 wt. .%, more preferably 0.5 to 1.25 wt.%, 0.5 to 1.15 wt.%, or 0.6 to 1.4 wt.%, 0.7 to 1.3 wt.%, or 0.8 to 1.1 wt.%.

Accordingly, Cr is present in Al-based powders and articles made from said powders in an amount of 0.1 to 1.4% by weight, preferably 0.2 to 1.4% by weight, preferably 0.3 to 1.4% by weight, more preferably 0.4 to 1.4% by weight, However, even more preferably it is present in an amount of 0.5 to 1.4% by weight.

For the same reason, Mn is present in amounts of 3 to 5.5% by weight in pre-alloyed Al-based powders and in articles made from such powders.

In general, it has been observed that within this concentration region of Mn, the alloy of the present invention responds to the content of Zr and Cr, but not outside. The specific co-dependence of Zr and Cr on Mn has been observed to result in a total concentration of Mn+Zr+Cr in a preferred embodiment of 3.4 to 8.5 wt.% in the Al-based powder and articles made from the powder.

In a preferred embodiment thereof, the total concentration of Mn+Zr+Cr is 3.7 to 8.0 wt.%, 4 to 7.5 wt.%, 4.25 to 7.25 wt.%, by weight in the Al-based powder and in articles made from said powder. %, or 4.5 to 7 wt.%.

Optionally, Mg may be present in up to 2% by weight of the pre-alloyed Al-based powder and in articles made from the powder. Mg will contribute to solution strengthening. In one embodiment, the pre-alloyed Al-based powder does not have Mg intentionally added, and is therefore considered another unavoidable impurity in the embodiment.

The presence of too much Fe and/or Si will cause the formation of low melting liquids in the final stages of the solidification process because these liquids can become trapped at grain boundaries and form cracks at the final shrinkage during solidification. Surprisingly, it was found that this phenomenon is independent of the individual contributions of Fe and Si, but only depends on the total concentration of these two elements, and that this phenomenon can be suppressed by limiting the total amount of Fe and Si (see Figure 4).

Therefore, an important feature appears to be maintaining the content of Fe and Si at a low level, at most 0.7% by weight of the elements in total, preferably at most 0.5% by weight, and more preferably at most 0.2% by weight in total.

The amount of other unavoidable impurities is at most 0.5% by weight, preferably at most 0.3% by weight of the Al-based powder and in articles made from the powder.

The Al-based powder according to the invention has a melting range of up to 300° C., where the melting range is defined as the difference between the onset of solidification temperature and the temperature when all the liquid has solidified.

However, to avoid cracking in the final stages of solidification, the drop in melt temperature close to complete solidification must be limited (see Figure 1). One practical way to evaluate the tendency of an alloy to form cracks during solidification is to use the solidification cracking susceptibility index.

The solidification cracking susceptibility index is defined according to equation (1):

where S is the solidification cracking susceptibility index, T is the temperature, and f _s is the mass fraction of solid.

The Al-based powder according to the invention has a solidification cracking susceptibility index of less than 20°C. The index is a measure of how the final melt solidifies, ranging from 80 to 100% solidified. The index measures the slope of temperature change over this range. A360.0 has a solidification cracking susceptibility index of 44°C, and Al 6160 has an index of 257°C. The lower the index, the less susceptible the alloy is to solidification cracking.

Al-alloyed powders can be manufactured by any rapid solidification process such as gas atomization, as well as by spray deposition, melt spinning, melt extraction, etc. without forming excessive precipitates, otherwise excessive precipitates may be produced by , these deposits will be detrimental to the manufacturing process, for example, by their tendency to block spray nozzles or otherwise obstruct the molten stream.

Powders with the intended composition are preferably prepared by melting a material with the desired composition and subjecting the melt to inert gas spraying to obtain a pre-alloyed spray powder. It is also possible to spray using air or water as the spray medium, but in the context of the present invention, if water is used, the oxygen content will be too high, and also the particle morphology may be damaged and the air and water atomized powder will be inert gas. This will have a negative impact on powder production as it tends to form more irregularly compared to powders formed by spraying.

The alloying elements of the powders of the present invention are also selected to prevent clogging of the spray nozzles during spraying. This problem has been avoided in some prior art powder compositions by using either pure elemental powders themselves or mixtures of different pre-alloy powders to obtain the desired chemical composition. A disadvantage of mixtures is that the different powders may separate during handling and processing, resulting in chemical non-uniformity of the printed part. This is avoided in the present invention because the aluminum powders of the present invention are simple to pre-alloy without loss of composition.

As is well known in the field of LPBF-printing, the size and shape of the powder particles used in printing are important to the ability to spread the powder in a uniform and homogeneous powder layer. It is known that workpiece cuts are identified at intervals of 5 to 150 μm.

Small spacing is advantageous for the melting process and the homogeneity of the powder bed, resulting in a more stable melting process with fewer defects such as porosity.

Typical particle size ranges are 15 to 45 μm or 20 to 53 μm, but depending on the LPBF equipment used and application requirements, any of the following sieve cuts may be used; 5-36, 10-45, 15-45, 20-53, 20-63, 45-90, 45-106, 53-106, 45-150, 53-150, 63-150, 75-150, 90- 150, 106-150 μm. A particular advantage of the Al-based powders of the present invention is that they can be readily manufactured in the size range desirable for LPBF-printing.

The stated particle size range means that at most 2% by weight of the powder have a particle size above the upper limit and at most 2% by weight of the powder have a particle size below the lower limit. Particle size was measured by laser diffraction according to ISO 13320-1 1999.

The shape (morphology) of the powder is also important in defining its diffusion behavior. The spherical shape results in more stable flow and diffusion behavior, resulting in fewer defects such as porosity and surface quality. Therefore, gas atomization is preferred when producing new Al-based powders because the metal powders produced by this method exhibit a spherical shape.

Al-based powders according to the present invention can function in several additive manufacturing processes, provided the solidification rate is sufficient. AM by electron beam melting and directed energy deposition are processes that meet these requirements, and the process is illustrated below using a laser powder bed fusion process.

Example

Example 1

An integrated computational materials engineering (ICME) strategy was applied to design the composition of Al-based powder. The alloy composition of the new alloy was optimized using CALPHAD (CALculation of Phase Diagram) type calculations.

CALPHAD calculations include a sufficient thermodynamic description of all possible phases in composition space for Al-alloys. Using this basis, predictions of the expected phases and their compositions are made with good accuracy. The results are valid for equilibrium conditions.

In materials prepared by LPBF, microstructures and phases are present as a result of rapid solidification. Equilibrium calculations are not relevant to the analysis of these microstructures. However, thermodynamic and phase relationships in composition space can be used to evaluate the tendency of phase formation by correlating a kinetic description of phase formation. To accomplish this, diffusion processes, nucleation and growth were considered along with thermodynamic explanations.

CALPHAD type calculations were used to predict the separation tendency of alloying elements in the final melt region as well as the solidification path. This separation should be limited to a crack-free solidified structure.

In the first step, a tentative composition was determined using calculations to obtain an alloy with a stable solidification path to avoid hot cracking.

The composition space for the alloy was then calculated and the maximum limits for alloying elements were determined. Composition limits were determined by calculating the phases that would form upon rapid solidification. Outside the limits indicated in Table 3, undesirable phases may form, causing problems in the spraying process or reducing mechanical properties.

By combining the calculation of phase equilibria with a kinetic description of the alloying system, the nucleation, growth and coarsening of the secondary hardening phase were simulated. By this method the alloy composition was adjusted to optimize precipitation of the secondary hardening phase during heat treatment.

CALPHAD calculations were coupled with a model that accounts for the different strengthening mechanisms, solution strengthening, precipitation strengthening and grain purification strengthening.

The main strengthening mechanism for these alloys was found to be solution strengthening. Precipitation hardening and grain refinement strengthening were also found to contribute to the strength of the alloy. The combination of these strengthening mechanisms determines the properties of the new alloy.

Table 3 below shows the results of calculations performed to determine alloying extent with respect to the formation of undesirable phases. To achieve solution strengthening, the highest possible amount of alloying elements was identified without forming unwanted phases.

Compositions marked with an "x" in Table 3 are acceptable with respect to the presence of undesirable phases, while compositions without any "x" are unacceptable with respect to alloying effects.

The formation of Cr-containing precipitates in the melt or solid phase will reduce the solution hardening contribution because they remove Cr from the solution. Precipitates formed at high temperatures in the melt will coarsen and will not provide a precipitation hardening effect.

The formation of Mn-containing precipitates in the melt or solid phase will reduce the solution hardening contribution. Large primary precipitates do not contribute to precipitation hardening. Too large a fraction of primary deposits will cause clogging of the spray nozzles and production stoppages.

Table 3

Example 2

Four different gas atomized Al-based powders, A, B, C, and D, with chemical compositions according to Table 4 were prepared using computer-derived alloy compositions. Several parts were printed using different printing parameters. In particular, laser power and laser speed varied over a wide range.

Large areas were found where the relative density of the manufactured article was greater than 98.5% (see Figure 2).

Further examination by optical microscopy revealed a single-phase structure without any cracks. Vickers hardness according to ISO 6507-1 - 2018, HV0.3, was measured in as-printed as well as heat-treated conditions, and the results are presented in Figure 3. Alloys A, C and D are age hardenable, and it can be seen that the hardness in particular for alloys C and D increases significantly to just below 130 HV0.3.

graph. 4

Effects of Chromium

From the experiments, it can be deduced that the effect of chromium is mainly solution strengthening, but importantly and surprisingly, chromium also affects the precipitation hardening sequence. This produces an alloy with improved properties and temperature resistance.

The as-printed samples are precipitation-free (verified experimentally), meaning that all precipitation is induced during the curing heat treatment. Therefore, the order of precipitation is influenced by composition.

Precipitation hardening heat treatment at 678K hardens the alloy by sequential precipitation of Al ₆ (Mn,Cr), Al ₁₂ (Mn,Cr), and Al ₃ Zr. The amount of chromium in the alloy will dictate the precipitation sequence and determine the amount of precipitates and their stability against coarsening. This means that the properties are critically dependent on the specific amount of Cr as detailed herein.

The advantage of the alloys of the invention comprising chromium compared to similar materials without chromium is higher strength, but importantly and surprisingly also increased temperature stability, which is an important parameter for industrial applications.

The effect of Cr is demonstrated in Figure 3. The cure curve shows:

● Alloy A (AlMnZr): Starts around 90 HV and peaks only around 105 to 110 HV.

● Alloy B (AlMnCr): Starts at approximately 98 HV and peaks at 648 K to 115 HV.

● Alloy C (AlMnCrZr - present invention): starts at -103 HV and peaks at 648 K to 140 HV.

● Alloy D (AlMnCrZr - invention): starts at -103 HV and peaks at 648 K to 140 HV.

This experiment clearly shows how chromium changed the cure rate. Alloy A hardens at 20 HV (90 -> 110 HV), while alloys C and D harden at 35 HV (103 -> 140 HV). Cure speed reflects the precipitation sequence, while as-printed hardness reflects the effect of solution cure. The effect of chromium on the alloy is a combination of both.

Here, a surprising synergistic effect is observed whereby the rearrangement of Mn and Cr between different precipitates slows down the coarsening rate and the hardening effect is maintained for a longer time, providing better temperature stability.

The curve shows that chromium has a significant effect on the onset of hardening. Alone, chromium will not provide an alloy with the desired properties. When chromium is combined with zirconium, its hardening effect is maintained for a longer time.

The alloy with only Mn and Zr but no Cr has significantly lower hardness. Higher contents of Mn and Zr will not improve this.

Example 3

The solidification behavior for Alloy D according to Example 2 and the known wrought alloys 6081, 7075 and A360.0 was simulated using the Scheil-Gulliver method and the CALPHAD method, as also used in Example 1.

Figure 1 shows the results of the simulation. Note that only the region close to complete solidification is included in the diagram, and the solid fraction ranges from 0.8 (80%) to 1 (100%). It can be seen that wrought alloys 6061 and 7075 exhibit slope curves, with 7075 in particular exhibiting a steep slope near the end of solidification. This behavior is typical for materials that form hot cracks at the end of solidification. Cast alloy A360.0 shows much better behavior with only a slight drop close to complete solidification. Alloy D of the present invention does not exhibit this drop at the end of solidification.

The individual solidification cracking susceptibility index, S, was determined for each alloy and the results are shown in Table 5 below.

Table 5

Example 4

From the detailed experiments above (Example 1 and Table 3), it was shown that, individually, Fe and Si could be present at up to 0.7% wt without negatively affecting the suitability of the Al-based powder of the invention for LPBF-printing. It turns out. This is advantageous because it reduces strain on the raw materials for use in forming the Al-based powders of the present invention, and less expensive recycled Al-based products can be used to form the Al-based powders of the present invention.

In numerical experiments as detailed above (Example 1), the effect of Fe and Si on the cracking susceptibility of the alloys of the present invention was determined when both of these impurities were present in the raw materials used to form the Al-based powders of the present invention. , were further investigated to determine whether limits exist for the content of Fe and Si.

As shown in Figure 4, the solidification temperature was found to be dependent on the total Fe and Si contents. The higher the total content of these elements, the greater the risk of forming a melt with a low solidification temperature that will cause problems with hot-cracking in LPBF-printed objects using the powders of the present invention.

Therefore, both Fe and Si individually can be present up to 0.7% wt without causing the presence of a melt with a solidification temperature that is too low, but their combined concentration should also not exceed 0.7% wt. From Figure 4, it follows that preferably the total amount of Fe and Si should not exceed 0.5%wt, and even more preferably should not exceed 0.2%wt.

Example 5

In further experiments with LBPR-printed objects reported herein, powder variants compared to the experiments under Example 2 were prepared by a nitrogen gas atomization process, and the produced powders fell into the nominal particle size range of 20 to 53 μm. It has been done.

Four categories of powder grades were designed with final chemical compositions as shown in Table 6. The particle size distribution was determined by laser diffraction according to ISO 13320-1 1999 using a Mastersizer 3000 from Malvern UK and the measurements were repeated five times (see Table 7 and Figure 5).

The stated particle size range means that at most 2% by weight of the powder have a particle size above the upper limit and at most 2% by weight of the powder have a particle size below the lower limit.

The experiments clearly demonstrated the suitability of the Al-based alloy of the present invention for forming Al-based powders with high uniformity.

Table 6: Alloy composition

Table 7: Particle size distribution

Example 6

LB-PBF processing and experimental design to establish overall density were tested in subsequent experiments.

Different variants of the powder from Example 5 were processed on an EOS M100 machine with a 40 μm spot size, 200 W (170 W nominal power) Yb-fiber laser. Samples were initially printed with a normal scan rotation of 67° maintaining 170 W power, 0.13 mm hatch distance, 1000 mm/s speed and 0.03 mm layer thickness as main processing input.

Powder samples were conditioned by a drying procedure at 353 K for 4 hours before every printing.

Samples were printed as 10 mm x 10 mm x 10 mm cubes designed to indicate the direction of gas flow in the chamber (shown in Figure 6). After printing, the sample was cut with a saw.

The processability of the material was further inferred based on a simple cubic factorial design to confirm the condition of interest with high relative density (approximately 99.5%), as shown in Figure 2.

This design of experiment (DOE) was performed using hatch distance and laser speed as initial variables compared to the main machining inputs to identify conditions for high-density machining.

Afterwards, a second DOE was performed to ensure that the values were experimentally acceptable by checking the processing parameter window and printing three additional samples for all points of interest.

Finally, a third DOE was performed on one of the quaternary alloys (alloy D) to fabricate a complete three-dimensional design with laser power as the third variable in the DOE.

Table 8 summarizes the range of processing parameters used for each alloy in the DOE.

Table 8: Summary of processing parameters for design of experiment (DOE)

The relative density of the sample was then determined from light microscopy of the cross section. The irradiated section was observed along three planes as shown in Figure 6.

To determine and retrieve porosity, results for all three planes were obtained and averaged for each sample to minimize orientation effects on the overall porosity values subsequently reported. Image analysis was performed using ImageJ software and sections of each sample with an area of approximately 30 to 50 mm ² per side.

A first DOE was performed to evaluate the printability of all four alloys with the goal of reaching >95% relative density.

Printing parameters were developed from appropriate settings for the EOS M290 machine and modified to suit processing on the EOSM1OO with compensation for lower power and smaller laser spot size.

Based on this, the cubic DOE was set with a fixed laser power of 170 W and a layer thickness of 0.03 mm, with a scan speed of 500 to 1500 mm/s and a hatch distance of 0.1 to 0.15 mm.

A total of 13 samples were printed with these DOEs for each alloy, and the results were measured as surface plots as shown in Figure 8.

A second DOE was then performed to confirm the existence of a target relative density for the above-mentioned parameters.

A total of 12 samples were printed and the resulting surface plots are shown in Figures 7a) and b).

According to this second DOE, multiple samples (60 to 75 samples) for all alloys are printed using parameter settings for maximum relative density (170 W power, 0.1 mm hatch distance and 1500 mm/s laser speed). It was established that the relative density far exceeds the target to be used.

Finally, by varying the laser power depending on the hatch distance and speed, a DOE was also designed with a similar concept containing 21 samples. Still, the range of processing parameters was kept conservative to expect only high density samples (>98% relative density). Surface plots generated from the third DOE are shown in Figures 9a)-b).

The final processing window was inferred after overlapping all results from the three DOEs, showing the processing parameter window resulting in relatively high density (>99%) according to the surface plots shown in Figures 9c)-d).

Example 7 - Microstructural evaluation and mechanical testing

All samples were prepared by cutting close to the center of all three sides (shown in Figure 6) and then mounting them in an epoxy-based thermoset called Polyfast (from Struers). The samples were then ground/polished according to standard Struers manufacturing for Al-alloys on a Struers TegraPol 31 machine. Samples were etched using standard Keller reagent to make the melt pool boundaries visible.

Example 8 - X-ray diffraction

X-ray diffraction (XRD) of the sample was scanned between 20° and 100°2θ with a step size of 0.007° and scan step time as 68.59 s, Cu source (Kα = 1.5406 Å) on finely ground samples (up to 2000 grit size) using a Bragg-BrentanoHD X-ray machine.

X-ray diffraction was performed on all four alloys (A, B, C, D) under as-printed conditions. This was performed on samples with high density (>99%).

In the XRD plot presented in Figure 10, it was observed that all four alloys showed similar information from the peaks. The Al peak was found to shift to a higher angle (20°) compared to the position expected for pure Al, which is attributed to the solute dissolved in solid solution. Small peaks observed for all four alloys were observed at 40.4° and 43.0° (see enlarged view in Figure 8b), which were also attributed to Al-Mn precipitates, which were later confirmed through SEM analysis.

Figure 11 represents an attempt to correlate the content of each alloying element in the alloy with the Al peak shift. This means that the effect of each element is described in the literature [T. Uesugi and K. Higashi, “First-principles studies on lattice constants and local lattice distortions in solid solution aluminum alloys,” Computational Materials Science, vol. 67, pp. 1-10, 2013], using a 'rule of mixtures' approach calculated by first principles calculations for binary Al alloys. The only modification to the calculations was to correct the starting point of pure-A1 to 4.0478 Å instead of 4.016 Å as assumed in Uesugi (PDF nr. 0040787), because this value is much higher according to the DIFFRAC.EVA software. This is because it has been shown to be trustworthy.

Therefore, the effect of each element was plotted against 4.016 Å, and the linear fit was extrapolated starting at 4.0478 Å instead of 4.016 Å (shown in Figure 10A). The values were then verified against the experimental value deduced from the strongest Al-peak at about 45°, as shown in Figure 10b). The experimental and calculated lattice parameter can be observed to be 0.01 Å.

Example 9 - Light Microscopy

Optical microscopy was performed on a ZEISS Axioscope 7 instrument with an automated scale that allows stitching of cross-sectional images up to 50 mm ² at 10X optical zoom.

In the as-spray state, the dendritic structure of the powder variant was clearly visible after etching with Keller's reagent. Figure 12 shows an optical microscope image for Alloy C illustrating the dendritic structures formed during the spraying process.

Low magnification optical microscopy images from all three cut sections of Alloy D were also taken as-printed after polishing to reveal a nominally fully dense structure with an average relative density of 99.56%. Representative images are shown in Figure 13 for all three cutting planes as previously defined.

When analyzing all samples, two types of common defects were recognized as observed for LPBF printed objects of aluminum alloy. The first one in Figures 14(a)-(b) was believed to represent keyhole porosity or gas porosity that occurs when the laser power is too high or the scan speed is too small; This means that the molten layer is the point of particle emission or evaporation in certain cases. The other in Figures 14(c)-(d) shows the lack of fusion type porosity that occurs when there is too low a laser power or too high a scan speed, resulting in melting the layers and bonding them to neighboring layers. The conditions are not sufficient.

Therefore, the observed porosity likely represents the maximum porosity under non-optimized LPBF-printing conditions.

Example 10 - Electron Microscopy

Microstructural evaluation was performed on a Leo Gemini 1550 SEM equipped with a field-emission gun. Depending on the type of imaging contrast required, imaging was performed with secondary electron (SE) and backscattered electron (BSE) receivers.

For compositional analysis, X-ray energy dispersive spectroscopy (EDS) for microanalysis was performed with a secondary electron acceptor coupled with INCA X-sight software.

Selected samples were analyzed using a Zeiss Gemini SEM 450 scanning electron microscope (SEM) with a field emission gun source for microanalysis. The microscope was fitted with a Bruker Quantax FlatQuad energy dispersive X-ray spectroscopy (EDX) detector, which allows elemental mapping of microstructures at sub-micron resolution.

The as-sprayed powder exhibited a pure dendritic structure for all alloys. Figure 15a illustrates a cross-sectional view of a particle of Alloy C. High-magnification imaging additionally revealed some fine nanometer-sized particles (<100 nm) inside the dendrites, as shown in Figure 15b. The chemical composition of these particles could not be determined because they were too small to be resolved using EDS in SEM.

After printing, the cubes were polished and lightly etched before characterization in an SEM in backscatter mode (BSE) to allow easy recognition of grain contrasts. Images along the front and top sides are shown in Figure 16, showing the apparent columnar structure along the build direction.

Upon further examination of the front surface, it was observed that two different categories of precipitates were formed in the alloy, characterized by their shape and location in the sample:

1. Sediment located at the boundary of the melt pool

2. Precipitates located at grain/sub-grain boundaries

The precipitates located at the melt pool boundary are shown in Figure 17, where it can be seen that these precipitates have two typical sizes: a spherical shape and a smaller one with an average size of <100 nm; The larger ones are 200 to 500 nm in size.

The second category of precipitates were those that formed in grain/subgrain (cell) systems. Their shape was rod-like with a length of 200-400 nm and a small diameter (<50 nm). The sediment is very clearly visible in Figure 18.

By itself, secondary phase particles aligned along grain/subgrain boundaries are assumed to be indicative of cellular solidification. Particles may form along these boundaries during solidification, resulting in a cellular solidification structure.

Figure 19 shows a summary based on EDS point scans of these three types of deposits. EDS has much greater diffusion for smaller particles due to the X-ray interaction volume generated by the electron beam of the SEM compared to the volume of sediment. However, to compensate for this, a larger number of point scans (>25) were performed at high resolution to obtain consistent results.

The results showed that the smaller melt pool boundary and subgrain boundary precipitates were richer in Mn than the matrix concentration, but could not indicate whether these would be either Al ₁₂ Mn (metastable) or Al ₆ Mn (stable). However, the larger melt pool boundary precipitates appear to be close to the Al ₆ Mn stoichiometry.

After illustrating the size, morphology and chemistry of the secondary phase particles as sprayed and as printed, the chemical composition of the matrix was evaluated using EDS chemical analysis in SEM. The matrix was generally observed to be supersaturated with elements in solid solution, with the exception of nanometer particles formed at melt pool boundaries or grain/subgrain boundaries. The EDS analysis of the matrices of all alloys in the sprayed (A-A) and printed (A-P) states is summarized in Table 9. In the alloy composition, the remainder is aluminum.

Table 9: Chemical composition via EDS for all four Al-alloys in sprayed (A-A) and printed (A-P) conditions.

Occasionally, a portion of the melt pool was found to have a refined grain structure. This effect is shown for alloy D in Figure 20b). A few Zr-containing particles with faceted or dendritic-like structures were also observed in the as-printed state (see Figure 21). Note that the size of these structures varies depending on the scale. The formation of large faceted grains is not clear, and the average composition of both precipitates is shown in Table 10, which shows that in the faceted Al-Zr structure seen in alloys C and D in Figure 21(a)(b). The apparent composition according to EDS is shown.

Table 10: Apparent sediment composition

Experimental considerations

Solidification cracking, which is common in LPBF-printing of certain Al-alloys, was addressed in this study. This problem was solved by ensuring that the alloy composition of the present invention contains elements that are not prone to separation during solidification.

This study introduces a new aluminum alloy family comprising two ternary and two quaternary variants. The alloy composition was designed to take advantage of the unique processing conditions provided by the rapid solidification and remelting that occurs during additive manufacturing when using laser-based powder bed fusion.

These alloy design principles follow the prediction of alloy solidification using ThermoCalc software. This enabled a simple route to completely prevent solidification cracking. The basis of alloy design is to develop aluminum alloys that can contain higher amounts of solute in supersaturated solid solution, which can then be used to achieve high strength through solid solution strengthening integrated with secondary phase precipitate strengthening.

These alloys have been shown to be resistant to hardening at up to 523 K under as-printed conditions and are therefore expected to be suitable for high temperature applications. Microhardness results showed that upon aging at 678 K, the average hardness moves from 105 HV at as-printed conditions to 130 HV at potential peak aging conditions.

Solidification cracks can appear to arise from shifting concentration gradients in the liquid as solidification occurs, thereby resulting in a higher solute content in the final liquid, resulting in higher volumetric contraction when such liquid solidifies. These cracks form between solidifying grains and may span multiple grains. This means that there is a minimum temperature gradient, thus ensuring a good amount of liquid metal available to bond along the two grain boundaries and prevent crack formation. According to the Scheil solidification curve, the solidification cracking susceptibility for the new Al-alloy developed in this regard was found to be significantly reduced. The experimental design supported this crack resistance, so all samples produced were shown to be crack-free in as-printed condition.

LPBF processing involves higher cooling rates due to the rapid melting and solidification of thin layers (usually 20 to 40 μm) of metal powder, resulting in cooling rates in the range of 103 to 105 K/s. This provides alloy designers with the opportunity to assume a higher supersaturation of elements in the Al-matrix during AM processing.

The alloying elements, and compositions, were selected based on studies performed on fast solidifying materials in which the benefits of transition metal elements into solid solutions of aluminum alloys were used to advantage.

EDS results show that almost all Mn and Cr are dissolved in the matrix, except for small precipitates formed during solidification (subgrain boundary precipitates) and remelting of the layers visible close to the melt pool boundary. This amount of precipitation is surprisingly small compared to other alloys known in the art.

Since the effect of remelting and separation of layers along the solidification front is always present during processing, the formation of secondary precipitates cannot be completely prevented in the as-printed state, and nanometric phases containing Mn, Cr were observed. In the art, segregation of Mn along grain boundaries has been shown experimentally by other authors in fast-solidifying nanocrystalline Al-Mn-based alloys, which also show that higher concentrations of Mn have 5-fold symmetry with the Al-matrix. This suggests that it leads to the formation of a 'quasicrystalline' phase, a metastable phase with a composition close to Al ₆ Mn.

It is suggested that the strength of the printed version of the alloy is a mix of solid solution strength and precipitate strengthening. Nanometric precipitates help restrict the movement of dislocations, and their relatively small size helps them become more efficient in strengthening the matrix.

As previously mentioned, much higher strengths can be achieved after aging by balancing the coarsening of grain/subgrain boundary precipitates and the nucleation/growth of new nanometric Ll ₂ Al ₃ Zr precipitates.

Since some of the nanometric precipitates containing Mn and Cr may be pseudocrystals or precursors of pseudocrystals, heat aging treatments were developed to control their formation and prevent them from forming the stable orthorhombic Al ₆ Mn phase. needs to be

Strengthening of this _new family of alloys by _Al . Advantageously, the as-printed version of the alloy is resistant to aging at 523 K for up to 24 hours.

Closing comment

Although the invention has been described in detail for purposes of illustration, such details are for that purpose only and it is understood that modifications may be made by those skilled in the art from a study of the drawings, disclosure, and appended claims in practicing the claimed subject matter. I understand.

It should be understood that the embodiments shown in the drawings are for illustrative purposes only and should not be construed as limiting the invention. Unless otherwise indicated, the drawings are intended to be read in conjunction with the specification (e.g., cross-hatching, arrangement of parts, proportions, scale, etc.) and are to be considered a part of the entire written description of the disclosure.

As used in the claims, the term "comprising" does not exclude other elements or steps. The indefinite article “a” or “an” used in the claims does not exclude plurality. A single processor or other unit may perform the functions of several means recited in the claims. Reference signs used in the claims should not be construed as limiting the scope.

Claims (12)

Pre-alloyed Al-based powders suitable for additive manufacturing consisting of the following total pre-alloyed Al-based powders by weight:
3 to 5.5% by weight Mn,
0.2 to 2% by weight Zr,
0.2 to 1.4% Cr by weight,
0 to 2% by weight Mg,
Fe and Si totaling up to 0.7% by weight,
0 as unavoidable impurities up to 0.7% by weight, and additional unavoidable impurities up to 0.5% by weight,
The rest of Al. 2. Pre-alloyed Al-based powder according to claim 1, wherein the Zr content is 0.3 to 1.8% by weight. 3. Pre-alloyed Al-based powder according to claim 1 or 2, wherein the Cr content is 0.3 to 1.4% by weight. 4. Pre-alloyed Al-based powder according to any one of claims 1 to 3, wherein the total content of Fe and Si is at most 0.2% by weight. 5. The pre-alloyed Al-based powder according to any one of claims 1 to 4, wherein the pre-alloyed Al-based powder does not contain intentionally added Mg. 6. Pre-alloyed Al-based powder according to any one of claims 1 to 5, having a solidification cracking susceptibility index S according to equation (1) of less than 20°C. 7. The particle size range of 10 to 53 μm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles are outside the specified particle size range. Pre-alloyed Al-based powder. A method of manufacturing an article by additive manufacturing, comprising:
- providing a pre-alloyed Al-based powder according to any one of claims 1 to 7;
- depositing a layer of said pre-alloyed Al-based powder to form a powder layer;
- heating the pre-alloyed Al-based powder according to a predetermined pattern by laser or electron-beam to fuse or melt the particles to be bonded;
- cooling the fused or molten particles;
- repeating the deposition, heating and cooling steps until an additively manufactured article is formed;
- Recovering the manufactured article
Method, including. 9. The method of claim 8, followed by heat treatment. 10. The method of claim 9, wherein the heat treatment comprises heating the manufactured article to a temperature of 150 to 450° C. for a period of at least 0.5 hours in an air atmosphere. An article formed by a melt fusion process of pre-alloyed Al-powder according to any one of claims 1 to 7. An article formed by a thermal fusion process of a pre-alloyed Al-powder according to claim 11, wherein the melt fusion process is according to any one of claims 8 to 10.